Silicon carbide possesses properties that should make it a superior semiconductor for applications that involve high temperature, high power, high radiation and/or high frequency (See Amorphous and Crystalline Silicon Carbide and Amorphous and Crystalline Silicon Carbide II). In addition, a variety of optical devices (e.g. light-emitting diodes, LEDs) can be fabricated from SiC and operated at high temperature. The properties that allow this superior performance are its large bandgap, excellent physical stability, high thermal conductivity, high electric breakdown field, and high saturated electron drift velocity. Semiconductor devices fabricated from SiC are capable of operating at temperatures above 600.degree. C. which is well above the capability of current commercial semiconductors. Also, based on theoretical considerations, SiC microwave devices should be far superior to currently available devices. The potential market for SiC devices is very large and could exceed billions of dollars annually.
The chief obstacle to the commercialization of SiC has been the lack of control over its crystal growth. Several properties of SiC contribute to this lack of control. First, it does not melt at reasonable pressures; it sublimes at temperatures above 1800.degree. C. Second, it grows in many different crystal structures, called polytypes.
The following defines some of the nomenclature used in this document. The SiC polytypes are formed by the stacking of double layers of Si and C atoms. Each double layer may be situated in one of three positions, designated as A, B, and C. The sequence of stacking determines the particular polytype and the stacking direction is called the crystal c-axis. There is one cubic polytype, with the zincblende structure and known as 3C or .beta. SiC. It has a three-layer repeat sequence ABC . . . , and so forth. All of the other polytypes are known as .alpha.-SiC and have either a hexagonal or rhombohedral structure. The hexagonal 6H-SiC polytype has the six layer repeat sequence ABCACB . . . , and so forth. For the .alpha.-SiC polytypes, the (0001) plane is known as the basal plane, and this plane is perpendicular to the c-axis. The (111) plane of 3C-SiC is equivalent to the basal plane of the .alpha.-SiC polytypes. In the discussions of this document, "basal plane" shall refer to either the (0001) plane for .alpha.-SiC, or the (111) plane of 3C-SiC. Also, "3C" shall be used for 3C-SiC, and "6H" shall be used for 6H-SiC. It is well known that a SiC surface, which is approximately parallel to the basal plane, is terminated with either Si atoms (the surface called the Si face) or terminated with C atoms (the surface called the C face). The term "vicinal (0001) wafer" shall refer to SiC wafers whose misorientation (i.e. tilt angle) is less than 6.degree. from the basal plane. The term "homoepitaxial" shall refer to epitaxial growth whereby the film and the substrate (wafer) are of the same polytype, and the term "heteroepitaxial" shall refer to growth whereby the film is of a different polytype than the substrate.
Since melt-growth techniques cannot be applied to SiC, vapor growth processes have been developed. An early SiC vapor-growth process, the Lely process, was based on the sublimation of polycrystalline SiC within a growth cavity and did produce rather pure crystals of various polytypes. Unfortunately, the crystals were too small, irregular in shape, and thus not suitable for commercial development. More recently, a high temperature sublimation process has been developed that does produce large single-crystal boules of 6H. Furthermore, polished wafers, more than 25 mm in diameter, can be produced from these boules. Commercial 6H-SiC devices are now being produced with these wafers.
Each of the various SiC polytypes has unique electrical and optical properties that can give it advantages over other polytypes in particular applications. For example, the 6H polytype has an energy bandgap of 2.9 eV and a hexagonal structure, whereas the 3C polytype has a lower bandgap, 2.2 eV, and a zincblende structure with a higher symmetry than 6H. These property differences lead to advantages for 6H in some applications (e.g. the wider bandgap--resulting in blue light emitting diodes, and possibly higher operating temperatures compared to 3C). On the other hand, the differences lead to advantages for 3C (higher electron mobility--leading to possibly higher frequency operation, and narrower bandgap--leading to lower forward voltage drop in some devices, compared to 6H). Thus it can be seen that it is desirable to develop devices based on 3C as well as 6H.
As of now, there is no existing method for producing single-crystal 3C boules. Hence, no 3C wafers are available. It is well known that single-crystal homoepitaxial 6H films can be grown in the lower temperature range 1400.degree. C. to 1550.degree. C. by chemical vapor deposition (CVD) on vicinal (0001) 6H wafers if the tilt angle is greater than about 1.5.degree.. Typically 3.degree. to 4.degree. is used If the tilt angle is less than about 1.degree. then a 3C film will be produced on the 6H wafer if prior art processes are used. However, the 3C film grown in this way generally has a high density of defects, including a defect known as double positioning boundaries (DPBs). This DPB defect can arise because of the change in stacking sequence of the 6H wafer (i.e. ABCACB . . . ) to that of the 3C (ABC . . . or ACB . . . ) film at the interface between the two polytypes. The difference between the two 3C sequences is a 60.degree. rotation about the &lt;111&gt; axis. If both of these sequences nucleate on the 6H substrate, DPBs will form at the boundary between domains differing by the 60.degree..
A theoretical crystal growth model, proposed by Matsunami and which has been used to explain the formation of the 3C and 6H polytypes on vicinal (0001) 6H substrates, is based on the density of atomic-scale steps on the growth surface. According to this model, 6H grows on 6H when the tilt angle is greater than about 1.degree. because terraces between steps are small and arriving molecules containing Si and C are able to migrate to steps where growth occurs. This growth is a lateral growth of the steps and reproduces the 6H substrate. For small tilt angles, say less than 1.degree., the terraces are larger and all molecules are not able to migrate to steps; instead, nucleation of 3C takes place on the terraces. Hence, for small tilt angles, 3C grows on 6H substrates. It has been reported that 3C films with reduced DPBs were grown on the as-grown face of 6H Lely-grown crystals, but in that case, it was assumed that the as-grown face was nearly atomically flat.
According to the model of Matsunami, nucleation of 3C on a given terrace would yield 3C nuclei with the same stacking sequence. Hence, no DPBs would be caused by the coalescence of such nuclei. In the case of polished wafers, the density of steps will always be relatively high, leading to a high density of DPBs. Thus, for the case of polished wafers, prior art teaches that a high density of DPBs is inevitable. There is generally a high density of stacking faults in the vicinity of DPBs indicating high stress in these regions. Also, films with high DPB density have a rough morphology. Such films are not suitable for device fabrication.
In the CVD growth of epitaxial SiC films on vicinal (0001) .alpha.-SiC wafers, a variety of pregrowth process have been used to prepare the polished surface for growth. The intent of these processes is to remove contamination and near-surface defects (from cutting and polishing the wafer) that contribute to poor quality films. Processes used to eliminate near-surface defects include molten-salt etching, oxidation followed by removal of the oxide with hydrofluoric acid, reactive ion etching, etc. prior to loading wafers into the CVD growth system. In situ processes (within the CVD growth system) include high temperature etching in H.sub.2 or HCl/H.sub.2 mixtures. All prior art use of these processes in CVD at temperatures less than 1600.degree. C. has produced the heteroepitaxial growth of 3C on vicinal (0001) 6H if the tilt angle was less than 1.degree.. In one case, Powell et al. used an HCl etch consisting 2 min at 1200.degree. C. In another case, Matsunami et al. used an HCl etch consisting of 10 min at 1500.degree. C. In both of these cases, 3C films were produced on vicinal (0001) 6H with tilt angles of less than 1.degree.. It must be emphasized at this point that prior-art teaches that pregrowth surface treatments of substrates used in CVD processes, can be effective in reducing defects in the resulting film. Prior art has not taught that surface treatments can be a significant factor in controlling the polytype of CVD-grown SiC films.
U.S. Pat. No. 5,248,385 discloses that surface treatments can be used to obtain homoepitaxial growth of SiC polytyptes on SiC substrates (e.g. 6H-SiC on 6H-SiC) with low tilt angles.
Copending application Ser. No. 718,315 filed Jun. 12, 1991, now abandoned discloses that growth of SiC crystal polytypes can be controlled through the use of appropriate surface treatments for growth surfaces with low tilt angles. This application also teaches that the location of the intentional nucleation site on the intended growth surface is important and should be located preferably at the highest atomic plane. However, it does not teach that the tilt angle at this preferred location (i.e. the highest atomic plane) should have any preferred value (e.g. zero tilt angle).
Additionally, in the abandoned patent application describes several methods for producing intentional heteroepitaxial growth at the desired location. specifies "a prescribed localized alteration of the crystal substrate surface." Other methods specify that the substrate surface can be altered by (1) scratching or indenting the surface with a diamond tool, (2) striking an electric arc to the surface, and (3) heating the surface with a laser, an electron or ion beam, or some other sort of energy beam. All of these techniques damage the surface in one way or another and the resulting heteroepitaxial nucleation is difficult to control. Another method specifies that intentional heteroepitaxial nucleation can be generated by "the introduction of a suitable impurity at the site to stimulate the growth of a desired polytype (e.g. 2H-SiC, 3C-SiC, etc.)." In the "DETAILED DESCRIPTION OF A PREFERRED EMBODIMENT", it is suggested that the vapor-liquid-solid (VLS) method of growth can be used to produce a single crystal whisker of a desired polytype. The 2H-SiC polytype is given as an example. But, this suggestion does not teach any procedure for carrying out the heteroepitaxial growth using the VLS technique. To those skilled in the art, it is known that VLS has been used to produce only the homoepitaxial growth of SiC whiskers and not heteroepitaxial growth of SiC whiskers on SiC substrates. This prior art does not provide any procedure for producing an intentionally grown heteroepitaxial "seed" crystal at the preferred site of heteroepitaxial nucleation. Thus, it is obvious to those skilled in the art that improvements need to be made in the procedure for achieving intentional heteroepitaxial nucleation at the preferred nucleation site.
Additionally, in abandoned patent application Ser. No. 718,315 filed Jun. 12, 1991, the methods contained therein describe that the heteroepitaxy be carried-out using a suitable silicon-containing source and a carbon-containing source. This prior art does not teach that the ratio of the silicon-containing source to the carbon-containing source plays any role in enhanced intentional heteroepitaxial nucleation. Also, it does not teach that altering any of the other growth conditions (e.g. temperature) can have any effect on the intentional heteroepitaxial nucleation.